ABSTRACT

This book chronicles the proceedings of the First International Symposium on Adhesion Aspects of Thin Films, held in Newark, New Jersey, October 28-29, 1999. Films and coatings are used for a variety of purposes a decorative, protective, functional, etc. a in a host of applications. Irrespective of the intended function or application of a film

chapter |1 pages

oy, is given as:

chapter |4 pages

A cknowledgements

chapter x|2 pages

-y

chapter |2 pages

Acknowledgements

chapter 2|9 pages

1. Interface stitching

chapter |40 pages

predicting the permissible external loading that a diamond-coated cutting tool can withstand without premature de-bonding. 3.1.6. Wear mechanisms. The failure of CVD diamond-coated inserts during machining can be in the form of flaking (interfacial failure) or abrasive wear (gradual cohesive failure) [22]. Ideally, a test of superb adhesion is when the diamond coating fully deteriorates by wear rather than flaking. Flaking will occur primarily due to poor adhesion between the diamond coating and the carbide substrate [6]. Therefore, flaking is clearly undesirable because the benefit of using a diamond coating is lost, except for the chip breaking assistance of faceted diamond crystals at the rake surface [29, 75]. If the adhesion strength of the CVD diamond coating is sufficient to withstand the machining stresses, then the abrasive action between the workpiece material and the diamond coating becomes the primary failure mechanism. Unless the CVD diamond coating is polished, a two-step wear mechanism is ex­ pected to occur. The first step is caused by the initial high surface roughness of the CVD diamond coating in which crack initiation occurs at the surface. The mecha­ nism that describes such behavior was proposed by Gunnars and Alahelisten [56]. They described a three-zone wear model as shown in Fig. 6. In this model, the role of residual stresses becomes significant in controlling crack propagation from the surface to the interface that could lead to interface failure (flaking). As outlined earlier, the high total compressive residual stress present in CVD diamond coatings on carbide inserts was assumed to be biaxial and oriented parallel to the interface. Wear starts to occur at the surface, which, because of geometry, allows stress to relax. A crack is more likely to initiate at protruding grains in zone I and propa­ gate preferentially along the (111) easy cleavage planes of diamond. The geometry at deeper depths, however, prevents the compressive residual stress from relaxing. Therefore, as the crack propagates deeper in the coating, it encounters higher com­ pressive stresses that cause the cracks to redirect their paths deviating from cleavage planes to a direction parallel to the interface in region II. The high compressive stress now causes cracks to propagate fast parallel to the interface resulting in a smooth surface in region III. Due to the smoother surface, fewer asperities will be present and it becomes harder to nucleate cracks.

chapter c|6 pages

= [ 1 +

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mechanics approach, similar values for the fracture energy are calculated; although the failure of the film under the indenter tip suggests a possible embrittlement of the ductile Al film in the presence of contaminants. The evidence of fracture within the carbon layer supports similar trends reported in whisker pull-out tests in which a carbon layer formation has been reported to decrease the fracture resistance [36]. Weak bonding has been reported for a range of interfacial carbon layer thicknesses. Brennan [37, 38] has cited an extensive fiber pull-out occurring in composites with a 10-40 nm layer of carbon present, attributed to debonding within the carbon layers. The effect on shear strength was recently reported by Dehm et al. [39] in which a decrease in the adhesion of copper to sapphire was observed by incorporating up to 110 nm thick carbon layers. Our current investigation shows that the debonding within a carbon layer occurs in much thinner layers than previously reported. The presence of a carbide suggests that

chapter 3|5 pages

2. Ceramic/interlayer coatings

chapter |3 pages

expansion with a heterodyne laser interferometer (laser probe). Demodulation is obtained with specific electronics. The magnitude and phase of the surface vibration are given with a second lock-in amplifier (lock-in amplifier 1) and stored in a microcomputer that also drives the scanning units. With this multi-acquisition microscope, the typical duration of an experiment in order to obtain a set of five low noise images is about 15 minutes. The resolution of the SThEM is given by the size at the photothermal source (radius of the optical beam: 5 /xm here). 4.1. Application to the study of thin films The first example concerns the observation of subsurface thin layers. In order to demonstrate the capacity for subsurface investigation we successively vapour deposited a 200 nm thick SiC>2 and 100 nm thick aluminium layers onto a polycrystalline nickel substrate (Fig. 8a). The bright strip on the right part of the image (Fig. 8b) reveals the presence of the subsurface SiC>2 layer which is optically invisible. This image has been obtained at 220 kHz modulation frequency of the excitation beam. The image contrast corresponds to about 25° phase shift. As the SThEM makes it possible to observe the subsurface we decided to use it for the detection of thin films delamination. We used a 1 /xm thick DLC film deposited on a steel substrate. Several lines of Vickers indentations were performed under an applied load of 4.5N. A different spacing (25 to 140 pim) between indentations has been taken for each line. The SEM and thermoelastic images of the indentations spaced 25 /xm are shown in Fig. 9. Due to the film delamination, an optically invisible bright area between the indentations (Fig. 9a) was observed by the SThEM at 100 kHz operating frequency (Fig. 9b). It is an indication of the excessive heating resulting from the film delamination. The latter is due to the tensile residual stresses which develop around each indentation. The bright area (film delamination) could not be detected both in the case of a single indentation or when the spacing between indentations was higher than 40 /xm. In the latter case

chapter |2 pages

is to identify these adhesional strains by using the radius of curvature and to find the stress in each material. We assume that the total strain (£tot)at any point of the system represented in Fig. 2, is: (Navier-Bemouilli’s hypothesis). Therefore, the effect of transverse shear (rxy = 0) is neglected, (ii) the radius of curvature is large compared with transverse dimensions [width (b) and thickness (h) of the three-layer system], leading to R\ b\, hh (iii) longitudinal elements of the beam are subjected only to simple tension or compression inducing stresses in the x direction, (iv) Young’s modulii of the coating having bulk properties, the interphase and the substrate have the same value in both tension and compression (flexural modulus). Based on these assumptions, final uni-axial residual stresses (a), in the x -direction of the three-layer system (bulk coating/interphase/substrate) are given by: of the zero deformation (y0)- Therefore, we consider two equilibrium conditions for the force (N) and the moment (M) for any cross section (area S) of the coating/interphase/ substrate system:

with £mech is the mechanical strain. Considering the geometry and the size of the three-layer systems studied, the beam theory can be used. For the one-dimensional approach, without lateral (width- wise) stresses, the following assumptions are made: (i) transverse sections of the beam are planar before, during, and after bending where: yo is the position where total strains are equal to zero (£tot = 0), y is the coordinate distance of any longitudinal fiber, R\ is the radius of curvature and E is Young’s modulus. To determine the distribution of residual stresses in the tri-layer system from equation (10) requires a knowledge of the radius of curvature (R\) and the position

chapter |11 pages

Coating thickness [mm]

chapter |7 pages

Dynamic mechanical thermal analysis (DMTA). Dynamic viscoelastic

Xt variation versus the coating thickness we 2.2.4. Nuclear magnetic resonance spectroscopy (NMR). For proton and carbon

chapter |17 pages

Floating part Precipitate

aluminum and titanium modified DGEBA are shown in Fig. 10. All spectra are

chapter 2|4 pages

2. Adhesion of the plasma-polymerized fluorocarbon films to silicon substrates The adhesion properties of the plasma-polymerized FC coatings were determined by using a test, already employed by Yasuda and Sharma [13] (see Fig. 1 and Table 1) in which the silicon substrates coated with plasma FC-films were boiled in a0.9% sodium chloride solution. The FC thin films produced in the processes 1 and 2 were lifted after a very short time (15 minutes). Coatings generated in process 3 were lifted after the second cycle of boiling. The films produced in processes 4 and 5 withstood the complete test procedure. The results are shown in Fig. 3. The poor adhesion of the polymerized films in the first two processes is due to the fact that these processes do not involve a plasma pre-treatment process. The difference between processes 1 and 3 is only in the plasma pre-treatment (process 1 does not contain the pre-treatment step of the silicon surface). The fluorocarbon films deposited by processes 4 and 5 have shown the best adhesion. These test results indicate that the plasma pre-treatment is very important and necessary for a good adhesion of the FC coatings to the silicon surfaces. 2.3. Patterning of FC films 2.3.1. Patterning through resist mask. The patterning of the FC films through a photoresist mask (conventional All resist AR-P351) was examined after deposition for process No. 5. Different coating parameters were investigated to improve the adhesion of the resist to the FC surface. The best adhesion results were obtained using the process parameters, shown in Table 3. Differences in the thickness uniformity of so-deposited resists were in a range below 5%. The samples were etched in a pure oxygen plasma in an RIE-system after the lithography steps (pre-bake, exposure, development, post-bake). A resolution of 2 /xm was obtained. A significant increase in the surface energy was not observed after resist stripping. The sessile contact angle of water was 103°. 2.3.2. Lift-off process for patterning thin plasma polymerized FC films. A lift-off process was also examined to pattern the thin FC films. The lithography steps were used before the plasma polymerization process was carried out (Fig. 2). A standard resist AR-P351 was coated directly onto the Si substrates. After all lithography

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